NANOSTRUCTURED MATERIALS

Intergranular corrosion resistance of nanostructured austeniticstainless steel

A. T. Krawczynska • M. Gloc • K. Lublinska

Received: 8 November 2012 / Accepted: 1 March 2013 / Published online: 12 March 2013

� The Author(s) 2013. This article is published with open access at Springerlink.com

Abstract Nanostructured metals and alloys possess very

high strength but exhibit limited plasticity. Enhancement of

the strength/ductility balance is of prime importance to

achieve wide industrial applications. However, post-

deformation heat treatment, which is usually used to

improve plasticity, can lead to a decrease in other proper-

ties. In the case of austenitic stainless steels, heat treatment

in the range from 480 to 815 �C can increase their sus-

ceptibility to intergranular corrosion. The aim of the work

reported in this paper was to determine if nanostructured

austenitic stainless steel is susceptible to intergranular

corrosion if heat treated for 1 h at 700 �C. Samples of

316LVM austenitic stainless steel were hydrostatically

extruded, in a multi-step process with the total true strain of

1.84 to produce a uniform microstructure consisting of

nanotwins. These nanotwins averaged 21 nm in width and

197 nm in length. Subsequent annealing at 700 �C pro-

duced a recrystallised structure of 68-nm-diameter nano-

grains. The heat treatment improved the ductility from 7.8

to 9.2 % while maintaining the ultimate tensile strength at

the high level of 1485 MPa. Corrosion tests were per-

formed in an aqueous solution consisting of 450 ml con-

centrated HNO3 and 9 g NaF/dm3 (according to ASTM

A262-77a). The evaluation of the corrosion resistance was

based on transmission and scanning electron microscopic

observation of the microstructure and chemical analyses.

The results revealed that both the as-received and

HE-processed samples are slightly susceptible to the

intergranular corrosion after annealing at 700 �C for 1 h.

Introduction

Nanostructured materials have attracted considerable

attention because of their higher strength, hardness and

wear resistance than the microstructural counterparts [1–4].

For example, precipitation-hardened nanostructured 2XXX

aluminium alloy after severe plastic deformation (SPD) by

hydrostatic extrusion (HE) possesses a strength similar to

that of conventional 7XXX alloys, which are considered to

be the strongest aluminium alloys [2]. Nanostructured

316LVM austenitic steel exhibits a lower friction coeffi-

cient under lubricated conditions than its microstructural

counterpart [4] and showed less intensive wear damage.

However, little is known about the other properties of

nanometals, including their corrosion resistance. Recent

research has shown that nanostructured 316 austenitic

stainless steel is more resistant to pit nucleation than the

conventional alloy [5]. It might be caused by the homog-

enization of the structure during HE. This favourable

behaviour of nanostructured austenitic stainless steel raises

a question about its resistance to intergranular corrosion.

This form of corrosion is particularly perilous for stainless

steel components operating in chemically aggressive

environments at high temperature. The major cause of the

susceptibility of austenitic steels to intergranular corrosion

is the formation of chromium carbide, Cr23C6, at grain

boundaries. This causes a depletion of chromium in the

areas surrounding these carbides. This possibility of this

form of corrosion must be taken into account when

applying post-deformation heat treatment to enhance the

ductility of nanometals [6–9], since the high strength is not

accompanied by high ductility in the case of nanometals.

Several methods have been proposed to avoid sensitising

stainless steel, and thereby eliminating the risk of inter-

granular corrosion, including a reduction of the carbon

A. T. Krawczynska (&) � M. Gloc � K. Lublinska

Faculty of Materials Science and Engineering,

Warsaw University of Technology, Woloska 141,

02-507 Warsaw, Poland

e-mail: akrawczynska@wp.pl

123

J Mater Sci (2013) 48:4517–4523

DOI 10.1007/s10853-013-7283-z

content, the addition of elements like titanium or niobium

and creating specially engineered grain boundaries [10–

12].

The aim of the reported work was to investigate whether

improving the mechanical properties by SPD and subse-

quently annealing the nanostructured 316LVM austenitic

stainless steel at 700 �C for 1 h affects its susceptibility to

intergranular corrosion.

Experimental

Sandvik Bioline 316LVM austenitic stainless steel, sup-

plied as cold-deformed 10-mm-diameter rods of the

chemical composition given in Table 1, was employed.

Billets cut from the rods were subjected to a multi-pass

HE process to achieve a final diameter of 4 mm, which

corresponds to a total true strain of 1.84. Samples of the as-

received and HE-processed material were annealed in air at

700 �C for 1 h. Before corrosion tests the samples were

ground and polished. Corrosion tests were performed in an

aqueous solution consisting of 450 ml concentrated HNO3

and 9 g NaF/dm3 (according to ASTM A262-77a). The

evaluation of corrosion resistance was based on micro-

structural observations by transmission and scanning

electron microscopes and chemical analyses.

Tensile testing was performed on the as-received, HE-

processed and annealed samples using a MTS Q Test/10 at

a speed of 10-3 m/s. The dimensions of the tensile test

samples are illustrated in Fig. 1.

Samples for the TEM examination were prepared by

mechanical polishing to produce a disc of thickness of

about 100 lm and further thinning, to obtain electron

transparency, was carried out by electropolishing. The

microstructures were examined using a JEOL JEM 1200

EX transmission electron microscope and a SEM 5500

scanning electron microscope. The microstructures were

quantified on the basis of the images obtained. The images

were analysed using computer-aided image analysis. Pre-

cipitates and etched dimples present in the samples after

the various treatments were analysed in terms of their size,

shape and chemical composition. The size of precipitates

and etched dimples were characterised as the equivalent

diameter, d2, defined as the diameter of a circle of area

equal to the surface area of the given precipitate or etched

dimple. To quantify the variation in size a variation coef-

ficient, CV (d2), defined as the ratio of the standard devi-

ation to the mean value, was applied. The shape was

described by the elongation parameter, which is the ratio of

the maximum to the equivalent diameter dmax/d2. The

phase analysis of selected carbides was undertaken by the

nanodiffraction mode using a high-resolution scanning

electron microscope HD2700.

Results

Microstructural observations before annealing

The microstructure of the as-received and HE-processed

samples consisted of deformation twins and shear bands.

The HE sample possessed significantly refined deformation

twins with an average length of 197 nm and width of

21 nm with a high density of dislocations, details of which

can be found elsewhere [13]. To achieve better character-

isation of the microstructure features like twins, dark-field

observations were performed, as illustrated in Fig. 2.

Tensile tests

The stress–strain curves of 316LVM austenitic stainless

steel in the as-received state, HE-processed, and HE-pro-

cessed and annealed at 700 �C/1 h are shown in Fig. 3 and

determined values are shown in Table 2.

It is evident that HE enhanced the strength. The yield

stress was increased from 839 to 1504 MPa and the ulti-

mate tensile strength from 1024 to 1793 MPa. However,

the uniform elongation dropped from 7.1 to 1.0 % and total

elongation from 24 to 7.8 %. After annealing the total

elongation increased from 7.8 to 9.2 % with only a slight

loss of strength. This indicated that significant improve-

ment in the strength–ductility balance can be achieved by

low-temperature annealing. It is well-known that this heat

treatment can increase the susceptibility of austenitic

stainless steels to intergranular corrosion. For this reason, it

is always necessary to perform microstructural observa-

tions to verify the presence of chromium carbides respon-

sible for intergranular corrosion.

Table 1 The chemical composition (wt%) of 316LVM steel

C Si Mn P S Cr Ni Mo Cu N

0.025 0.6 1.7 0.025 0.003 17.5 13.5 2.8 0.1 \0.1

Fig. 1 Dimensions of samples for tensile tests

4518 J Mater Sci (2013) 48:4517–4523

123

Fig. 2 a The microstructure of the as-received sample in the bright

field with the diffraction pattern in the orientation [011] of the matrix

and [0-1-1] of twins; b, c microstructures in the dark field of the

matrix and twins; d the microstructure of the HE-processed sample in

the bright field and the corresponding diffraction pattern with the

circled spot used in the dark field; e the microstructure of twins in the

dark field

J Mater Sci (2013) 48:4517–4523 4519

123

Microstructure after annealing at 700 �C for 1 h

The grain size and shape of the microstructures produced

by annealing at 700 �C for 1 h have been described pre-

viously [13]. It should be noted that nanocarbides were

evident in both the as-received and HE-processed samples,

as shown in Fig. 4. The nanocarbides are formed at the

intersections of shear bands and at the boundaries of the

twins and grains. In the as-received samples carbides

appear much more sparsely than in the HE-processed

samples. Therefore, only their average equivalent diameter,

39 ± 15 nm, was determined, which is comparable to the

average equivalent diameter of 33 ± 12 nm of the carbides

in the HE-processed samples. The carbides are elongated

which is expressed in their parameter of elongation (1.40).

Furthermore, their coefficient of variation is only 0.36,

which suggests that they are relatively homogenous in size.

Chemical analyses of the carbides in the as-received and

HE-processed samples after annealing at 700 �C/1 h were

undertaken. Results for a few chosen carbide particles are

presented in Fig. 5 and Tables 3 and 4. The chemical

analysis shows that precipitates rich in Mo and also Cr are

present in the as-received and HE-processed samples.

Cr-rich precipitates are particularly hazardous as their

formation reduces the chromium content in the area close

to the grain boundaries which induces intergranular cor-

rosion. The carbide particles in the grain boundaries are

strongly cathodic to the adjoining chromium depleted

zones. In the HE-processed sample, phase analysis was

carried out and an example of the diffraction pattern

obtained for a selected carbide particle is shown in Fig. 6.

Examination of the pattern gave clear evidence of the

presence of Cr23C6 carbide in the annealed structure.

Corrosion behaviour of annealed samples

Before commencing the corrosion tests the surface is

totally free of dimples or scratches as indicated in Fig. 7a,

b.The surfaces of samples after the corrosion tests, on

which etched dimples are visible, are shown in Fig. 7c–f.

The dimples were formed at the shear bands and at both

grain and twin boundaries. These dimples could be the

effect of corrosion of the areas surrounding the Cr- and

Mo-rich precipitates resulting from the reduced chromium

0

500

1000

1500

2000

0 5 10 15 20 25

Strain [%]

Str

ess

[MP

a]AS-RECEIVEDHE HE 700C/1h

Fig. 3 Stress–strain curves of austenitic stainless steel in the

as-received state, HE-processed, and HE-processed and annealed at

700 �C/1 h

Table 2 Mean values and standard deviation of ultimate tensile strength

YS UTS Au At

E (MPa) SD (MPa) E (MPa) SD (MPa) E (%) SD (%) E (%) SD (%)

As-received 839 24 1024 12 7.1 0.8 24.0 2.2

HE 1504 31 1793 21 1.0 0 7.8 0.1

HE ? 700 �C/1 h 1357 11 1485 8 1.0 0.3 9.2 3.6

UTS; uniform elongation, Au and total elongation, At in the as-received state, HE-processed, and HE-processed and annealed at 700 �C/1 h

Fig. 4 Microstructures of austenitic stainless steel after annealing a the as-received and b HE-processed samples

4520 J Mater Sci (2013) 48:4517–4523

123

content of areas. The etched dimples are evenly distributed

on the surface of the HE-processed samples. However,

examination at high magnification revealed that the etched

dimples did not form lines at grain boundaries. It is highly

probable that in both samples intergranular corrosion was

at the initial stage. The parameters of the size and shape of

the etched dimples and their volume fraction in the as-

received and HE-processed samples after annealing for 1 h

at 700 �C are shown in Table 5.

The average size of etched dimples in the as-received

samples was shown to be 0.96 lm and was three times

larger than those in the HE-processed samples. The etched

dimples in the as-received samples were of submicrometer

size, whereas in the HE-processed samples the etched

dimples were of both nano and submicrometer size. The

differences in size between the samples were well-characterised

by the coefficient of variation, which is much higher in the

HE-processed than in the as-received samples. On both

types of sample the perimeter of the etched dimples was

almost circular. In HE-processed samples the volume

fraction of etched dimples on the HE material of 9.0 % was

much greater than the 1.9 % exhibited by the as-received

material.

Discussion

It was shown that the simple technique of annealing

nanostructured 316LVM stainless steel for 1 h at 700 �C

increased the elongation to fracture from 7.8 to 9.2 %

whilst maintaining a high UTS of 1485 MPa.

However, the results indicate that the possibility of

intergranular corrosion is an important issue that must be

considered when contemplating post-deformation anneal-

ing of nanostructured stainless steels. Annealing for 1 h at

700 �C created more nanocarbide particles in HE-pro-

cessed materials that in the as-received material which

consequently produced a greater volume fraction of etched

dimples during the corrosion tests. This suggests that the

formation of a nanostructure promotes carbide formation

during annealing. However, the average equivalent diam-

eter of the dimples was smaller in the HE-processed sam-

ples and no etched lines were visible at the grain

boundaries in either condition which implies that the for-

mation of the nanostructure does not affect the rate of

corrosion.

It should be noted that sensitised steels might be more

prone to other types of corrosion, such as pitting or crevice

corrosion. For this reason, a novel approach should be

taken to improve the ductility of nanometals, especially

nanostructured steels. It has recently been proposed that an

unconventional process, annealing under high pressure,

achieves outstanding results [14]. Various conditions of

temperature and pressure have been analysed and, to date,

the best combination of ductility and strength, a ductility of

Fig. 5 Microstructures of austenitic stainless steel after annealing at 700 �C/h a, b the as-received c HE-processed samples; the carbides for

which chemical composition was determined are indicated

Table 3 Chemical composition of carbides (wt%) in the as-received

samples after annealing

Position Point Si Cr Mn Fe Ni Mo

Precipitate 1 (Fig. 5a) 1.97 16.09 1.28 44.55 8.78 27.33

Precipitate 2 (Fig. 5a) 2.95 13.42 1.00 29.79 4.49 48.35

Precipitate 3 (Fig. 5a) 2.59 15.56 0.78 40.34 7.29 33.43

Matrix 4 (Fig. 5a) 2.81 18.05 1.46 59.80 13.98 3.90

Precipitate 1 (Fig. 5b) 1.55 20.32 0.74 38.08 4.16 35.15

Matrix 2 (Fig. 5b) 2.55 18.17 1.93 59.38 14.10 3.87

Table 4 Chemical composition of carbides (wt%) in HE-processed

samples after annealing

Position Point Si Cr Mn Fe Ni Mo

Precipitate 1 (Fig. 5c) 3.02 11.64 0.30 27.53 3.35 54.15

Precipitate 2 (Fig. 5c) 1.78 28.13 0.66 36.69 4.80 27.94

Precipitate 3 (Fig. 5c) 2.83 11.47 0.36 27.09 3.98 54.27

Precipitate 4 (Fig. 5c) 3.09 12.32 0.30 26.60 3.39 54.30

Precipitate 5 (Fig. 5c) 2.77 12.67 0.40 27.96 4.58 51.63

Matrix 6 (Fig. 5c) 2.12 18.03 1.34 59.84 13.59 5.08

J Mater Sci (2013) 48:4517–4523 4521

123

24.4 % and strength of 1247 MPa, was obtained for HE-

processed 316LVM austenitic stainless steel by a heat

treatment consisting annealing at 900 �C for 10 min under

a pressure of 6 GPa. This treatment produced a

microstructure consisting of nanograins of the average

equivalent diameter of 130 nm. It was also very significant

that no carbide particles were created by this annealing

process. It is possible that annealing at the high pressure

Fig. 6 High-resolution picture of a selected carbide particle and diffraction patterns from the carbide (orientation [-111]) and matrix

(orientation [001])

Fig. 7 Surfaces of annealed samples before corrosion tests; a as-received; b HE-processed; surfaces of annealed samples after corrosion tests; c,

d as-received; e, f HE-processed

Table 5 Parameters of the size

and shape of etched dimples in

the annealed samples after the

corrosion tests

Sample Equivalent diameter d2 a Volume fraction (%)

of etched dimplesE (lm) SD (lm) CV

As-received ? 700 �C/1 h 0.96 0.31 0.32 1.21 1.9

HE-processed ? 700 �C/1 h 0.28 0.44 1.59 1.11 9.0

4522 J Mater Sci (2013) 48:4517–4523

123

retards the diffusion processes which lead to recovery,

recrystallization, grain growth and the formation of pre-

cipitates [15–19].

Conclusions

Nanostructured austenitic steel 316LVM was produced by

HE and subjected to post-deformation annealing. Tensile

tests showed that after annealing for 1 h at 700 �C the

material exhibited an enhanced ductility of 9.2 % and

retained a high tensile strength of 1485 MPa. However, the

heat treatment sensitised the steel and promoted inter-

granular corrosion. The research indicates that the possi-

bility of intergranular corrosion must be taken into account

when considering post-deformation annealing of austenitic

stainless steels.

Acknowledgements This work was carried out within a NANO-

MET Project financed by the European Funds for Regional Devel-

opment (Contract No. POIG.01.03.01-00-015/08). The advice and

assistance given by Prof. M. Lewandowska is much appreciated.

Open Access This article is distributed under the terms of the

Creative Commons Attribution License which permits any use, dis-

tribution, and reproduction in any medium, provided the original

author(s) and the source are credited.

References

1. Zhu YT, Langdon TG (2004) J Microsc 56:58

2. Lewandowska M, Kurzydlowski KJ (2008) J Mater Sci 43:7299.

doi:10.1007/s10853-008-2810-z

3. Garbacz H, Lewandowska M, Pachla W, Kurzydlowski KJ

(2006) J Microsc 223:272

4. Budniak J, Lewandowska M, Pachla W, Kulczyk M, Kurzyd-

lowski KJ (2006) Solid State Phenom 114:57

5. Pisarek M, Kedzierzawski P, Plocinski T, Janik-Czachor M,

Kurzydlowski KJ (2008) Mater Charact 59:1292

6. Dobatkin SV, Rybal’chenko OV, Raab GI (2007) Mater Sci Eng

A 463:41

7. Ma E (2006) J Microsc 58:49

8. Valiev RZ, Sergueeva AV, Mukherjee AK (2003) Scripta Mater

49(7):669

9. Wang H, Shuro I, Umemoto M, Kuo H, Todaka Y (2012) Mat Sci

Eng A 556:906

10. Korostlelev AB, Abramov VY, Belous VN (1996) J Nucl Mater

233–237:1361

11. Terrada M, Saiki M, Costa I, Padilha A (2006) J Nucl Mater

358:40

12. Jones R, Randle V (2010) Mater Sci Eng A 527:4275

13. Krawczynska AT, Lewandowska M, Kurzydłowski KJ (2008)

Solid State Phenom 140:173

14. Krawczynska AT, Brynk T, Rosinski M, Michalski A, Gierlotka

S, Grzanka E, Stelmakh S, Palosz B, Lewandowska M,

Kurzydlowski KJ (2012) Improving mechanical properties of

nanometals by annealing. In: Proceedings of the 33rd risoe

international symposium on materials science: nanometals—

status and perspective, 279–286

15. Tanner LE, Radcliffe SV (1962) Acta Metall Mater 10:1161

16. Syrenko AF, Klinishev GP, Khoi VT (1973) J Mater Sci 8:765.

doi:10.1007/BF00553726

17. Kuhlein W, Stuwe HP (1988) Acta Metall Mater 36:3055

18. Lojkowski W (1988) J Phys-Paris 49:545

19. Sursaeva V, Protasova S, Lojkowski W, Jun J (1999) Texture

Mictrostruct 32:175

J Mater Sci (2013) 48:4517–4523 4523

123